Ultra-High Strength Martensitic Alloy

ABSTRACT

An age hardenable martensitic steel alloy is disclosed. The alloy has the following composition in weight percent. 
                                       C   0.30-0.36                           Mn   0.05   max.         Si   0.05   max.         P   0.01   max.         S   0.0010   max.                       Cr   1.30-3.2          Ni   10.0-13.0         Mo    1.0-2.70         Co   13.8-17.4                           Ti   0.02   max.         Al   0.005   max.         Ce   0.030   max.         La   0.010   max.                                                          
The balance is iron and the usual impurities. The composition of this alloy is balanced to provide a unique combination of very high strength, together with good toughness, ductility, and resistance to fatigue.

FIELD OF THE INVENTION

The present invention relates to an age hardenable martensitic steel alloy, and in particular, to such an alloy which provides a unique combination of very high strength together with good toughness, ductility, and fatigue resistance.

BACKGROUND OF THE INVENTION

An age-hardenable martensitic steel alloy is known that provides a combination of high strength and good fracture toughness. The alloy is sold under the trademark AERMET® 310 and it has found wide utility in structural components for the aerospace industry, in armor for both air and land craft, and in machine tool components. The AERMET 310 alloy is capable of providing an ultimate tensile strength of about 2137 MPa (310 ksi) in combination with a K_(lc) fracture toughness of about 65.9 MPa√{square root over (m)} (60 ksi√{square root over (in)}).

The weight and size of structural components are critical design variables for the aerospace industry. This is also true in the automotive industry, particularly in the field of high performance racing cars. Therefore, aerospace design engineers and automotive engineers continually search for ways to reduce component size, and hence weight, without giving up important mechanical properties, particularly mechanical strength, toughness, and ductility. Because of this ongoing demand for materials that permit the use of reduced-weight structural components, it would be desirable to have a steel alloy that provides even higher strength than the AERMET 310 alloy. However, it is well known that the toughness and ductility of steel are typically inversely related to the strength property. Therefore, it is important that any such alloy provide the higher strength property without a significant loss in the toughness and ductility properties.

The function of certain structural parts for high performance automotive applications, such as springs, subjects such parts to high frequency, repetitive action, over long periods of time. Thus, a critical design property for the material used in such components is the resistance to failure from fatigue. Accordingly, for usefulness in the high performance auto racing, good fatigue resistance is needed in addition to the aforementioned combination of high strength, toughness, and ductility.

SUMMARY OF THE INVENTION

The combination of properties that are desired for the fields of use described above are realized to a significant degree by the steel alloy according to the present invention. The alloy according to this invention is an age hardenable martensitic steel alloy that provides significantly higher strength than the known alloy, while maintaining acceptable levels of toughness and ductility relative to the known alloy. In particular, the alloy of the present invention is capable of providing an ultimate tensile strength (UTS) of at least about 2344 MPa (340 ksi) with ductility and overall toughness at least similar to the AERMET 310 alloy. In addition, the alloy of this invention provides excellent fatigue resistance.

The alloy according to this invention is an age hardenable martensitic steel alloy having the following broad, intermediate, and preferred weight percent compositions.

Broad Intermediate Preferred C 0.30-0.36 0.32-0.35 0.32-0.34 Mn 0.05 max. 0.02 max. 0.01 max. Si 0.05 max. 0.05 max. 0.03 max. P 0.01 max. 0.005 max. 0.003 max. S 0.0010 max. 0.0010 max. 0.0005 max. Cr 1.30-3.2  2.0-2.5 2.20-2.30 Ni 10.0-13.0 11.0-13.0 11.5-12.5 Mo  1.0-2.70 1.5-2.2 1.80-1.90 Co 13.8-17.4 15.0-16.0 15.4-15.6 Ti 0.02 max. 0.015 max. 0.01 max. Al 0.005 max. 0.003 max. 0.001 max. Ce 0.030 max. 0.020 max. 0.010 max. La 0.010 max. 0.010 max. 0.005 max.

The balance of the alloy is essentially iron except for the usual impurities found in commercial grades of such steels and minor amounts of additional elements which may vary from a few thousandths of a percent up to larger amounts that do not objectionably detract from the desired combination of properties provided by this alloy.

The foregoing tabulation is provided as a convenient summary and is not intended to restrict the lower and upper values of the ranges of the individual elements of the alloy of this invention for use in combination with each other, or to restrict the ranges of the elements for use solely in combination with each other. Thus, one or more of the element ranges of the broad composition can be used with one or more of the other ranges for the remaining elements in the preferred composition. In addition, a minimum or maximum for an element of one preferred embodiment can be used with the maximum or minimum for that element from another preferred embodiment.

Here and throughout this specification, the term “percent” or the symbol “%” means percent by weight, unless otherwise indicated.

DETAILED DESCRIPTION

The alloy according to the present invention contains at least about 0.30% and preferably at least about 0.32% carbon. Carbon contributes to the good strength and hardness capability of the alloy primarily by combining with other elements, such as chromium and molybdenum, to form M₂C carbides during an age-hardening heat treatment. Too much carbon adversely affects fracture toughness, room temperature Charpy V-notch (CVN) impact toughness, and stress corrosion cracking resistance. Accordingly, carbon is limited to not more than about 0.36%, better yet to not more than about 0.35%, and preferably to not more than about 0.34% in this alloy.

Cobalt contributes to the very high strength provided by this alloy and benefits the age hardening of the alloy by promoting heterogeneous nucleation sites for the M₂C carbides. In addition, the contribution of cobalt to the very high strength provided by the alloy, is less detrimental to the toughness of the alloy than the addition of carbon. Accordingly, the alloy contains at least about 13.8%, better yet at least about 15.0%, cobalt. Preferably, at least about 15.4% cobalt is present in the alloy. Because cobalt is an expensive element, the benefit obtained from cobalt may not justify using very large amounts of it in this alloy. Therefore, cobalt is restricted to not more than about 17.4%, better yet to not more than about 16.0%, and preferably to not more than about 15.6%.

Carbon and cobalt are controlled in the alloy of the present invention to benefit the superior combination of very high strength and high toughness that is a characteristic of the alloy. I have observed that increasing the ratio of cobalt to carbon (Co/C) promotes increased toughness and a better combination of strength and toughness in this alloy. Further, increasing the Co/C ratio benefits the notch toughness of the alloy. Accordingly, cobalt and carbon are controlled in this alloy such that the ratio Co/C is at least about 43 and preferably at least about 52. However, the benefits from a high Co/C ratio are offset by the high cost of producing an alloy having a Co/C ratio that is too high. Therefore, the Co/C ratio is restricted to not more than about 100 and preferably to not more than about 75.

Chromium contributes to the good strength and hardness capability of this alloy by combining with carbon to form M₂C carbides during the age-hardening process. Therefore, at least about 1.30%, better yet at least about 2.0%, and preferably at least about 2.20% chromium is present in the alloy. Too much chromium increases the sensitivity of the alloy to overaging which can result in reduced strength. In addition, too much chromium results in increased precipitation of carbides at the grain boundaries, which adversely affects the alloy's toughness and ductility. Accordingly, chromium is limited to not more than about 3.20%, better yet to not more than about 2.50%, and preferably to not more than about 2.30% in this alloy.

Molybdenum, like chromium, is present in this alloy because it contributes to the very high strength and the hardness capability of this alloy by combining with carbon to form M₂C carbides during the age-hardening process. Additionally, molybdenum reduces the sensitivity of the alloy to overaging and benefits stress corrosion cracking resistance. Therefore, at least about 1.00%, better yet at least about 1.50%, and preferably at least about 1.80% molybdenum is present in the alloy. Too much molybdenum increases the risk of undesirable grain boundary carbide precipitation, which as noted above, may result in reduced toughness and ductility. Therefore, molybdenum is restricted to not more than about 2.70%, better yet to not more than about 2.2%, and preferably to not more than about 1.90%.

At least about 10.0%, better yet at least about 11.0%, and preferably at least about 11.5% nickel is present in the alloy to benefit hardenability and to reduce the alloy's sensitivity to quenching rate so that acceptable CVN toughness is readily obtainable. Nickel also benefits the stress corrosion cracking resistance and the K_(lc) fracture toughness. Too much nickel promotes an increased sensitivity to overaging. Therefore, nickel is restricted in the alloy to not more than about 13.0% and preferably to not more than about 12.5%.

Other elements can be present in the alloy in amounts which do not detract from the desired properties. Not more than about 0.05% and better yet not more than about 0.02% manganese is present because manganese adversely affects the fracture toughness of the alloy. Preferably, manganese is restricted to not more than about 0.01%. Also, up to about 0.05% silicon, up to about 0.005% aluminum, and up to about 0.02% titanium can be present as residuals from small additions for deoxidizing the alloy during melting. Preferably, silicon is restricted to not more than about 0.03%, aluminum is restricted to not more than about 0.003%, and titanium is restricted to not more than about 0.015%.

Small but effective amounts of elements that provide sulfide shape control are present in the alloy to benefit the fracture toughness property by combining with any sulfur present in the alloy. Such elements are effective to form sulfide inclusions that do not adversely affect the fracture toughness property. A similar effect is described in U.S. Pat. No. 5,268,044, which is incorporated herein by reference. In one embodiment of the present invention, the alloy contains up to about 0.030% cerium and up to about 0.010% lanthanum. Preferably, the alloy contains not more than about 0.020%, better yet not more than about 0.010% cerium, and not more than about 0.005% lanthanum.

The balance of the alloy is essentially iron except for the usual impurities found in commercial grades of alloys intended for similar service or use. The levels of such elements are controlled to avoid adversely affecting the desired properties. For example, phosphorous is restricted to not more than about 0.01% and preferably to not more than about 0.005% because of its embrittling effect on the alloy. Sulfur, although inevitably present, is restricted to not more than about 0.0010% and preferably to not more than about 0.0005%, because sulfur adversely affects the fracture toughness of the alloy.

The alloy of the present invention is readily melted using known vacuum melting techniques. For best results, a multiple melting practice is preferred and ultra-clean starting materials, such as electrolytic iron, are preferably used for charging the melting furnace. The preferred practice is to melt a heat by vacuum induction melting (VIM) and cast the heat in the form of an electrode. The electrode is then refined by vacuum arc remelting (VAR) into one or more ingots for further processing.

The preferred method of introducing cerium and lanthanum into this alloy is through the addition of high purity grades of cerium and lanthanum during VIM, prior to casting the VAR ingot electrode. Rare earth alloys such as NiLa may also be used. Effective amounts of cerium and lanthanum are present when the ratio of cerium to sulfur (Ce/S) in the VIM melt is at least about 4. When the Ce/S ratio is more than about 20, the retained amounts of the rare earths present in the as-cast VAR electrode ingot may adversely affect the hot workability and tensile ductility of the alloy. Preferably, the Ce/S ratio at VIM is at least about 8, and not more than about 10. In another embodiment of this alloy, a small but effective amount of one or more of calcium, magnesium, yttrium, or other sulfur-gettering element, is present in the alloy in place of a portion or all of the cerium and lanthanum to provide the beneficial sulfide shape control.

Prior to VAR, the electrode ingots are preferably stress relieved at about 677° C. (1250° F.) for 4-16 hours and air cooled. After VAR, the ingot is preferably homogenized at about 1177-1232° C. (2150-2250° F.) for about 6-24 hours.

The alloy can be hot worked from about 1232° C. (2250° F.) to about 816° C. (1500° F.). The preferred hot working practice is to forge an ingot from about 1177-1232° C. (2150-2250° F.) to obtain at least about a 30% reduction in cross-sectional area. The ingot is then reheated to about 982° C. (1800° F.) and further forged to obtain at least another 30% reduction in cross-sectional area.

Heat treating to obtain the desired combination of properties includes a solution treatment (austenitizing) and quenching, a deep chill treatment, and then an age hardening heat treatment. The alloy is austenitized by heating it at about 843-982° C. (1550-1800° F.) for about 1 hour plus about 5 minutes per inch of thickness, followed by quenching. The quench rate is preferably rapid enough to cool the alloy from the austenitizing temperature to about 66° C. (150° F.) in not more than about 2 hours. The preferred quenching technique will depend on the cross-sectional size of the manufactured part. However, the hardenability of this alloy is good enough to permit air cooling, vermiculite cooling, or inert gas quenching in a vacuum furnace, as well as oil quenching. After the austenitizing and quenching treatment, the alloy is preferably cold treated as by deep chilling at about −196° C. (−320° F.) for about 0.5-1 hour and then warmed in air. Age hardening of this alloy is preferably conducted by heating the alloy at about 454-510° C. (850-950° F.) for about 5 hours followed by cooling in air. The age-hardened alloy is preferably given an additional cold treatment at the same conditions specified above.

The alloy of the present invention is useful in a wide range of applications. The very high strength and good fracture toughness of the alloy makes it useful for structural components for aircraft and machine tool components. The alloy of this invention is also useful for automotive components including, but not limited to, structural members, drive shafts, springs, and crankshafts.

WORKING EXAMPLES

In order to demonstrate the novel combination of properties provided by the alloy according to the present invention, two experimental 400 lb. (181.4 kg) heats, Heats 1 and 2 were melted and processed into specimens for mechanical testing. The weight percent compositions of the experimental heats are set forth in Table 1 below. Also shown in Table 1 are the weight percent chemistries of four production heats of the AERMET 310 alloy (Heats A, B, C, and D) which were tested for comparison.

TABLE 1 HEAT HEAT HEAT HEAT HEAT HEAT ELMT 1 2 A B C D C 0.328 0.336 0.257 0.256 0.255 0.254 Mn <0.01 <0.01 0.02 0.03 0.02 0.02 Si <0.01 <0.01 0.02 0.03 0.02 0.02 P <0.005 <0.005 0.002 0.002 0.002 0.002 S <0.0005 <0.0005 0.0005 0.0008 0.0012 0.0005 Cr 2.24 2.22 2.46 2.45 2.45 2.46 Ni 11.01 12.07 11.08 11.02 11.04 11.10 Mo 1.82 1.81 1.45 1.45 1.46 1.44 Co 15.70 15.56 15.20 15.18 15.22 15.18 Ti 0.010 0.010 0.010 0.008 0.010 0.007 Al 0.015 0.011 0.004 0.003 0.004 0.003 Nb <0.001 <0.001 0.004 0.004 0.002 0.002 N <10 ppm <10 ppm 0.0010 0.0010 0.0010 0.0010 O <10 ppm <10 ppm 0.0010 0.0010 0.0010 0.0010 Ce 0.004 0.004 0.004 0.006 0.005 0.004 La 0.003 0.003 0.001 0.001 0.002 0.001 The balance of each heat is iron and the usual impurities.

Heats 1 and 2 were processed into forged bars measuring about 1.5 in. by 4.5 in. (3.81 cm by 11.4 cm). Duplicate specimens for tensile testing, Charpy V-notch testing, and fracture toughness testing were prepared from the forged bars of each of Heats 1, 2 and A-D. All of the test specimens were heat treated to provide maximum tensile strength. For Heats 1 and 2, the specimens were austenitized at 1775° F. (968.3° C.) for one hour, and then cooled in air. The specimens were then given a deep chill treatment at −320° F. (−196° C.) followed by warming in air. The specimens were age hardened at 900° F. (482.2° C.) for 5 hours, and then cooled in air. For comparative Heats A-D, the specimens were austenitized at 1675° F. (912.8° C.) for one hour, and then cooled in air. The specimens were then given a deep chill treatment at −100° F. (−73.3° C.) and then warmed in air. The specimens were age hardened at 875° F. (468.3° C.) for 6 hours, and then air cooled. Set forth in Table 2 below are the results of the mechanical testing of specimens including the tensile strength (UTS) and yield strength (YS), both in ksi (MPa), the percent elongation (% El.), the reduction in area (% R.A.), the Charpy V-notch impact strength (CVN), and the fracture toughness (K_(lc)) in ksi Vm (MPa Vm). CVN impact testing was performed in accordance with ASTM Standard Test E23. Fracture toughness testing was preformed in accordance with ASTM Standard Test E399.

An objective of the alloy according to the present invention is to maximize both strength and toughness. The relevant strength parameter is the ultimate tensile strength. Toughness, however, can be measured in numerous ways. Mechanical engineers often use a toughness measure which is an approximation of the area under the stress-strain curve. This measure allows them to design a part to “bend before breaking”. Parts using ultra-high strength alloys often are designed with toughness measures that take into account stress concentrations. The two most common tests to measure the effects of stress concentrations on toughness are the Charpy V-notch Impact Test and the Fracture Toughness test.

In order to provide an indication of the overall toughness of the alloy according to this invention, three toughness parameters were considered: the area under the tensile stress-strain curve, the CVN impact energy, and the fracture toughness (K_(lc)). Those three measures are combined into one parameter, a Toughness Index. The Toughness Index is the geometric mean of the three normalized toughness measures and is calculated as follows:

${{Toughness}\mspace{14mu} {Index}} = \sqrt[3]{\begin{matrix} {\left\lbrack {\left( {\left( {{Elong}.} \right) \times {\left( {Y.S.{+ {U.T.S.}}} \right) \div 2}} \right) \div 50} \right\rbrack \times} \\ {\left\lbrack {{CVN} \times 3} \right\rbrack \times \left\lbrack K_{Ic} \right\rbrack} \end{matrix}}$

Two of the toughness measures were “normalized” such that each of their values are within a 0-100 scale. Normalization was used so that the Toughness Index did not overly weight one toughness measure relative to another. In the foregoing equation, the area under the stress-strain curve ((Elong.)×(Y.S.+U.T.S.)/2) is normalized by dividing it by 50 and the CVN Impact Energy value is multiplied by 3. The fracture toughness value is used without normalizing. The Toughness Index values calculated for each of the experimental heats and two of the comparative heats are also presented in Table 2.

TABLE 2 TOUGH- % % NESS HEAT Y.S. U.T.S. el. R.A. CVN K_(Ic) INDEX 1 314.3 349.7 10.3 53.1  8.8 31.1 38.3 (2167) (2411) (11.9) 354.2 10.4 49.6 11.0 31.6 41.7 (2442) (14.9) 2 307.9 349.0 10.0 45.8 13.8 31.5 44.1 (2123) (2406) (18.7) 312.5 348.7 9.6 50.9 11.0 33.5 41.2 (2155) (2404) (14.9) A 290.3 335.0 9.4 40.6 12.0 (1) (1)  (2002) (2310) (16.3) 296.0 335.9 8.4 39.3 12.5 (2041) (2316) (16.9) B 285.8 332.3 6.0 20.9 20.0 40.7 44.9 (1971) (2291) (27.1) 289.5 331.5 9.2 38.4 25.5 42.5 48.3 (1996) (2286) (34.6) C 296.1 333.5 9.3 51.8 (1)  41.3 (1)  (2042) (2299) 287.2 330.4 11.3 54.7 40.7 (1980) (2278) D 287.0 335.4 9.3 33.8  7.3 40.6 37.2 (1979) (2313)  (9.9) 284.6 335.4 9.3 22.2  6.2 34.0 33.2 (1962) (2313)  (8.4) (1) Insufficient material to test.

The data presented in Table 2 show that Examples 1 and 2 of the alloy according to the present invention provide significantly higher tensile strength than any of the comparative Heats A, B, C, and D. The data, particularly the Toughness Index, also make clear that the toughness and ductility of Examples 1 and 2 are at least as good as the known alloy represented by the comparative heats. Thus, the combination of strength, ductility, and toughness provided by the alloy according to this invention is superior to the combination of those properties provided by the heats of the known alloy.

In order to demonstrate the good fatigue resistance of the alloy according to this invention, fatigue testing was performed on specimens of Example 1 and a further comparative heat, Heat E, having the following weight percent composition. Heat E is a production heat of the AERMET 310 alloy discussed above and was prepared similarly to the other comparative heats.

Heat E C 0.258 Mn 0.01 Si 0.02 P 0.002 S <0.0005 Cr 2.46 Ni 11.1 Mo 1.45 Co 15.23 Al 0.004 Ti 0.010 N <0.0010 O <0.0010 Ce* 0.006 La* 0.002 *At VIM The balance is iron and the usual impurities.

Set forth in Table 3 below are the results of the R.R. Moore Rotating Beam Fatigue Test for each of the alloys, including the number of cycles to failure (Cycles) and the applied stress (Stress) in ksi (MPa).

TABLE 3 Heat E Heat 1 Cycles Stress Cycles Stress 8.00 × 10⁴ 200 (1379) 1.00 × 10⁵ 200 (1379) 8.20 × 10⁴ 200 (1379) 1.83 × 10⁵ 200 (1379) 1.12 × 10⁶ 180 (1241) 6.39 × 10⁵ 180 (1241) 1.04 × 10⁶ 180 (1241) 6.99 × 10⁵ 180 (1241) 3.01 × 10⁶ 160 (1103) 7.06 × 10⁵ 160 (1103) 4.86 × 10⁶ 160 (1103) 3.01 × 10⁶ 150 (1034) 4.94 × 10⁶ 150 (1034) 6.70 × 10⁶ 160 (1103) 9.75 × 10⁶ 150 (1034) 6.84 × 10⁶ 150 (1034) 2.29 × 10⁷ 140 (965)  1.46 × 10⁷ 140 (965)  3.36 × 10⁷ 140 (965)  1.83 × 10⁷ 140 (965) 

The data set forth in Table 3 show that the fatigue resistance of Example 1 of the alloy according to the present invention is at least as good as the fatigue resistance of the known alloy.

The terms and expressions which have been employed in the foregoing disclosure are used as terms of description and not of limitation. There is no intention in the use of such terms or expressions to exclude any equivalents of the elements or features described or any portions thereof. It is recognized, however, that various modifications are possible within the scope of the invention as described and claimed. 

1. An age hardenable martensitic steel alloy, consisting essentially of, in weight percent, about C 0.32-0.36 Mn 0.05 max. Si 0.05 max. P 0.005 max. S 0.0010 max. Cr 1.30-2.30 Ni 11.5-13.0 Mo  1.8-2.70 Co 15.4-17.4 Ti 0.02 max. Al 0.005 max. Ce 0.030 max. La 0.010 max.

the balance being iron and usual impurities.
 2. (canceled)
 3. An age hardenable martensitic steel alloy as set forth in claim 1 that contains at least about 2.0 weight percent chromium.
 4. (canceled)
 5. An age hardenable martensitic steel alloy as set forth in claim 1 that contains not more than about 16.0 weight percent cobalt.
 6. (canceled)
 7. (canceled)
 8. An age hardenable martensitic steel alloy as set forth in claim 1 that contains not more than about 2.2 weight percent molybdenum.
 9. (canceled)
 10. An age hardenable martensitic steel alloy, consisting essentially of, in weight percent, about C 0.32-0.35 Mn 0.05 max. Si 0.05 max. P 0.005 max. S 0.0010 max. Cr 2.0-2.3 Ni 11.5-13.0 Mo 1.8-2.2 Co 15.0-16.0 Ti 0.015 max. Al 0.003 max. Ce 0.020 max. La 0.010 max.

the balance being iron and usual impurities.
 11. (canceled)
 12. An age hardenable martensitic steel alloy as set forth in claim 10 that contains at least about 2.20 weight percent chromium.
 13. (canceled)
 14. (canceled)
 15. (canceled)
 16. (canceled)
 17. An age hardenable martensitic steel alloy as set forth in claim 10 that contains not more than about 2.0 weight percent molybdenum.
 18. An age hardenable martensitic steel alloy as set forth in claim 10 that contains not more than about 0.34 weight percent carbon.
 19. An age hardenable martensitic steel alloy, consisting essentially of, in weight percent, about C 0.32-0.34 Mn 0.04 max. Si 0.03 max. P 0.003 max. S 0.0005 max. Cr 2.20-2.30 Ni 11.5-12.5 Mo 1.80-1.90 Co 15.4-16.0 Ti 0.01 max. Al 0.001 max. Ce 0.010 max. La 0.005 max.

the balance being iron and usual impurities.
 20. An age hardenable martensitic steel alloy as set forth in claim 1, 10, or 19 comprising an element selected from the group consisting of calcium, magnesium, yttrium, and combinations thereof, in place of at least a portion of the cerium and lanthanum.
 21. An article of manufacture formed from a martensitic steel alloy having a weight percent composition consisting essentially of about C 0.32-0.36 Mn 0.05 max. Si 0.05 max. P 0.005 max. S 0.0010 max. Cr 1.30-2.30 Ni 11.5-13.0 Mo  1.8-2.70 Co 15.4-17.4 Ti 0.02 max. Al 0.005 max. Ce 0.030 max. La 0.010 max.

and the balance of said alloy is iron and usual impurities, wherein said article provides a tensile strength of at least about 345 ksi (2344 MPa).
 22. An article of manufacture as claimed in claim 21 wherein said article is in the age hardened condition.
 23. An article of manufacture as claimed in claim 21 which contains more than 1.8% molybdenum.
 24. An article of manufacture as claimed in claim 21 wherein the ratio Ce/S is greater than
 4. 25. An article of manufacture as claimed in claim 24 wherein the ratio Ce/S is not greater about
 20. 